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In situ monitoring of corrosion processes within the bulk of AlMgSi alloys using X-ray microtomography
Corrosion Science 50 (2008) 3455–3466
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In situ monitoring of corrosion processes within the bulk of AlMgSi alloys using X-ray microtomography Fabian Eckermann a,b,*, Thomas Suter a,*, Peter J. Uggowitzer b, Andreas Afseth c, Alison J. Davenport d, Brian J. Connolly d, Magnus Hurlen Larsen e, Francesco De Carlo f, Patrik Schmutz a a
Laboratory for Corrosion and Materials Integrity, Empa, Swiss Federal Laboratories for Materials Testing and Research, Uberlandstrasse 129, 8600 Dubendorf, Switzerland Laboratory of Metal Physics and Technology, Department of Materials, ETH Zurich, 8093 Zurich, Switzerland c Alcan Inc., Voreppe 38341, France d University of Birmingham, Metallurgy and Materials, Edgbaston, Birmingham, B15 2TT, UK e Department of Materials Science and Engineering, Norwegian University of Science and Technology, N-7491 Trondheim, Norway f Advanced Photon Source, Argonne National Laboratory, Argonne IL, United States b
a r t i c l e
i n f o
Article history: Received 29 April 2008 Accepted 2 September 2008 Available online 23 September 2008 Keywords: A. Aluminium B. X-ray microtomography C. Acid corrosion C. Exfoliation corrosion
a b s t r a c t Susceptibility to localized corrosion of AlMgSi (AA6016 and AA6111) alloys in certain aggressive environments is high. In this study, synchrotron X-ray microtomography was used to monitor non-destructively corrosion processes within bulk materials. In the selected aggressive solutions, surface–deformed layers showed high localized corrosion susceptibility, but are more stable than the bulk of the alloy during corrosion propagation. In addition, exfoliation-like attack was observed as a transition from intergranular attack. This directed corrosion is not related to grain or crystallographic structure. Finally, intermetallic particles dissolution inside the materials after contact with aggressive solution of the intergranular corrosion path was evidenced. Ó 2008 Elsevier Ltd. All rights reserved.
1. Introduction X-ray microtomography is a non-invasive technique used to image three-dimensional (3D) bulk objects. Visualization of defects, dimensional inspection and local characterization are possible in the interior of a material [1]. The spatial resolution of the method can vary from several mm to the 100 nm scale. Use of higher spatial resolution limits the observed volume and acquisition times [2–7]. Measurement times also range from several hours to a few seconds depending on the source, detector, material and specimen dimensions. Over the past five to ten years, X-ray microtomography has become an increasingly accepted technique for qualitative and quantitative characterization of the 3D internal structures of materials. A very new field of tomography research using this capability is the study of corrosion processes. Whereas an initial in situ tomography experiment was performed in 2004 involving the corrosion of sandstone under triaxial compression and chemical atmosphere
* Corresponding authors. Address: Laboratory for Corrosion and Materials Integrity, Empa, Swiss Federal Laboratories for Materials Testing and Research, Uberlandstrasse 129, 8600 Dubendorf, Switzerland (F. Eckermann). Tel.: +41 44 823 4921; fax: +41 44 823 4015 (T. Suter). E-mail addresses: [email protected], [email protected] (F. Eckermann), [email protected] (T. Suter). 0010-938X/$ - see front matter Ó 2008 Elsevier Ltd. All rights reserved. doi:10.1016/j.corsci.2008.09.015
with a resolution of 0.1225 mm3 per voxel (=3D pixel) [8], studies of corrosion processes of metals have only started recently [9–11]. In corrosion studies it is of crucial importance to understand the changes occurring within the bulk of a material when it is exposed to an aggressive aqueous medium. In many localized corrosion sites of attack most damage takes place beneath the surface. However, most models describing localized corrosion have been deduced from ‘‘electro”chemical activity measured on the surface of a material or via traditional 2D, ex situ destructive analysis of metallographic cross-sections. From conventional methods it is difficult to obtain reliable information on complex corrosion propagation morphologies and on the interaction of microstructure with localized corrosion or corrosion cracking. The use of conventional methods (e.g. polishing) also enhances the risk of introducing artefacts during sample preparation. Grain fallout during polishing is one specific example. In the first publication concerning in situ characterization of corrosion for a metal, Marrow et al. [9] observed stress corrosion cracking development in stainless steel. Corrosion processes could be quantified and a 3D model was established. Recently Davenport et al. [10] used synchrotron smonochromatic X-rays (energy resolution of DE/E of 1.4 10 4) to correlate the yttrium distribution in a magnesium alloy with the corrosion path. In the first published work on in situ aluminium corrosion investigated using X-ray microtomography, Connolly et al. [11] demonstrated some of
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the corrosion phenomena occurring on an AA2024 alloy, such as intergranular corrosion path propagation and the dissolution of the matrix around intermetallics within AA2024. In this work three phenomena related to the corrosion of Al– Mg–Si alloys are investigated using X-ray microtomography: The role of intermetallic particles, the behaviour of a surface–deformed layer and the occurrence of exfoliation-like attack (ELA). 1.1. The role of the primary intermetallics in the corrosion process Considerable work has been carried out characterizing the cathodic trenching of intermetallic particles [12–14], often following initial dealloying reactions [15–21]. The role of certain intermetallic particles in pit initiation is also well established [17,18,22–26]. However, very few observations have been made of the behaviour of particles in the interior of materials, where aggressive solutions can build up in localized corrosion sites. 1.2. The behaviour of the surface–deformed and activated layer A surface–deformed layer is produced during any shearing event which takes place on the surface of aluminium, i.e. grinding, rolling or machining. This surface layer can induce filiform corrosion susceptibility depending on the formation of second phase particles and a fine grain structure [27]. The second phase particles are formed during heat treatments after the deformation process [27–30]. In the cited work this layer is of submicrometer size. However concepts presented for this surface–deformed layer might be transferred to surface–deformed layers of much larger thickness as present in this study. The stability of the surface–deformed layer after initiation of first localized attack has not so far been addressed in detail in literature. 1.3. Exfoliation-like attack (ELA) Exfoliation is a corrosion process assumed to occur in highly directional microstructures. A consensus on its mechanism has not yet been reached, but it has been attributed variously to intermetallics arranged in planes, to mechanisms similar to those for intergranular corrosion and stress corrosion cracking (in the 2xxx series), and also to segregation of alloying elements coming from the ingot [31,32]. Frankel stated for the AA2024 and AA7xxx alloys using a combination of different characterization techniques such as X-ray radiography, foil penetration technique and potentiodynamic polarization measurements that exfoliation follows intergranular corrosion and that the intergranular corrosion process moves along elongated grains faster in certain directions [33]. The kinetics of intergranular corrosion also depends on the grain aspect ratio [34]. Kelly claimed that exfoliation is triggered by stresses generated by corrosion products and follows an intergranular corrosion path [35]. Others have attributed exfoliation corrosion to segregation processes. The directed microstructure in the final rolled sheet is, for example, a remnant of the ingot microstructure [31,36,37]. None of the above-mentioned corrosion phenomena have so far been monitored and discussed for AlMgSi alloys in the context of X-ray microtomography investigations. This method has the advantage of monitoring of corrosion attack and tracking of corrosion paths as a function of time during corrosion processes which occur upon aggressive solution contact. It also makes it possible to characterize the morphology of corrosion propagation. In this way, not-yet-clarified changes in the behaviour of surface–deformed layers and intermetallics during corrosion propagation can be investigated. New information improving the understanding of the role of intermetallics, surface–deformed layers, corro-
sion morphology and indirectly about the solution chemistry can be gained. 2. Experimental X-ray microtomography experiments were performed at the Argonne National Laboratories, IL, USA, Advanced Photon Source beam line 2-BM [38]. Acquisition was carried out at a 17 keV Xray energy with a double Si 1 1 1 crystal monochromator. Fast acquisition of the radiographs was performed at 20 keV with a double multilayer monochromator. The detector CCD chip was read out with 1024 1024 pixels, giving a final practical resolution of about 3 lm. Seven hundred and twenty 2D radiographs were taken for the subsequent reconstruction procedure. The settings used for fast acquisition of the data lowered the acquisition time from 1 h to 10 min. For image processing, ImageJ software [39,40] was used. To measure corroded volumes, the tomograms were converted into binary images and a greyscale threshold was defined to allow separation of corroded and not-attacked regions. The set-up is designed so that a pin can be fully immersed in solution during tomographic measurements, whereby the dimensions of the pin are adjusted to allow short acquisition times (see Fig. 1). Samples for X-ray tomography experiments were first machined to 3 mm diameter rods, and then the upper 2 mm was turned down to 0.5 mm diameter pins. This thin part was analysed with X-ray tomography and is positioned in the middle of the sheet. The samples were lacquered with Stopping Off Lacquer to cover all but the to the X-ray beam exposed part of the pin. For samples investigated in situ, a silicone rubber tube with 3 mm internal diameter was slipped over the outside of the 3 mm rod to form a cell around the pin, and the cell was subsequently filled with solution. For the electron backscattered diffraction measurements (EBSD), a LEO GEMINI 1530 field emission gun SEM was used. The data were acquired at an angle of 70° with respect to the surface and 15 kV electron beam voltage. The aperture of the electron beam optics was opened to 120 lm. For corrosion tests performed on the sheet material, the surface was ground with 4000 SiC paper before being exposed to the solution.
Fig. 1. Sketch of exposure cell: a pin of 500 lm diameter was fully immersed in the solution confined in silicone rubber tubing. The field of view marks the area where the beam penetrated the specimen.
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2.1. Materials: AA6016 and AA6111 alloys The composition of aluminium alloys 6016 and 6111 is given in Table 1. The pins for the X-ray tomography experiments were cut from the center of 3 mm-thick aluminium sheets. The machined layer (surface–deformed) was not removed prior to the corrosion experiments. All samples used in this study were heat treated to 540 °C for 15 min (solution heat treatment) after machining and then water-quenched. Subsequent ageing treatment took place for 45 min at 180 °C, followed by water-quenching to achieve high intergranular corrosion susceptibility. For the EBSD mappings performed after corrosion experiments, the pin surfaces were polished to 0.05 lm with a vibromed polishing system. Fig. 2 shows the microstructure of the aluminium alloys used. Considering the cross-section through the AA6016 sheet (Fig. 2a), the microstructure obtained was very heterogeneous. Large grains in mm size dominate in the outer region; the center of the sheet consists of very small grains in the 10 lm-range. A similar heterogeneity was observed by Svenningsen et al. [41] for extruded AlMgSi, where large grains of over 1mm in size dominated the surface of the sheet, while in the center region fine grains were present. The nomenclature used for labelling the orientations is visualized in Fig. 3. No significant difference between the longitudinal (L) plane and longitudinal–transverse (LT) plane can be found in the grain structure. Within the AA6016 pins, the mentioned heterogeneous grain structure (fine and coarse grains) was present because the pins were cut out of the middle of the sheet material (shown in Fig. 3). In alloy AA6111 a different grain structure can be identified (Fig. 2b). The grains are slightly directed (aspect ratio of 3.6), and average a thickness of about 120 lm. Grain structure did not differ in size and shape in the LT or L plane.
Fig. 3. Schematic sketch showing the position of the pin in the sheet and the nomenclature used.
To investigate the corrosion processes of the AA6016 and AA6111 alloys, three types of corrosion experiments in particular were performed: (1) in situ with samples immersed in 0.7 M HCl; (2) in situ in 2.5 M HCl; and (3) ex situ after 0.7 M, 2.5 M and 5 M HCl solution exposure. The immersion times and solutions are listed in Table 2. AA6111 was immersed in 0.7 M HCl solution for up to 7 h and monitored in situ with 1 scan (X-ray tomography measurement) per hour. In situ experiments with specimens immersed in 2.5 M HCl were monitored with 1 scan every 15 min. Ex situ experiments were investigated after exposure to solution. The solution was refreshed after each scan to minimize unwanted effects as a change in the solution pH. Table 2 Overview of the experiments presented in this paper Experiment type
AA6016 (immersion time [h])
AA6111 (immersion time [h])
In situ
–
7
2
1
0.75 7 45 0.25 2 0.25
– 7 10 – 1 –
Ex situ Table 1 Composition of AA6016 and AA6111 measured by optical emission spectroscopy Wt%
Mg
Si
Fe
Mn
Cu
Cr
AA6016 AA6111
0.35 0.61
1.05 0.80
0.19 0.26
0.08 0.21
0.07 0.70
0.01 0.02
Ex situ Ex situ
0.7 M HCl 2.5 M HCl 0.7 M HCl 2.5 M HCl 5 M HCl
Type of material, immersion time and solution type are listed.
Fig. 2. Optical image of the alloys’ microstructure after Barker’s etching: (a) For AA6016 sheet material a significant difference in the grain structure through the sheet thickness is observed. (b) For AA6111 oriented the same way, homogeneous grain size distribution is visible.
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3. Results 3.1. Overview of the observed attack In this section, the newly-revealed phenomena of slow-dissolving surface–deformed layer and the transition from intergranular corrosion (IGC) to exfoliation-like attack (ELA) are illustrated. In Fig. 4, an ex situ SEM image of a corroded AA6016 pin used in the tomography experiments is shown, providing an overview of the corrosion damage observed upon immersion in aggressive HCl electrolyte. The pin (Fig. 4) was immersed for 45 h in 0.7 M HCl. The previously mentioned heterogeneous grain structure can be seen, as the pin is not attacked homogeneously. The preferred re-
Fig. 4. SEM image of a pin with heavy traces of corrosion. AA6016 immersed for 45 h in 0.7 M HCl.
gion of attack in the middle of the pin corresponds to the smallgrained section (Fig. 2a). However, the extent and morphology of the corrosion attack inside the pin is not predictable from such surface observation. In X-ray microtomography experiments the corrosion attack morphology within pins is visualized. Fig. 5 shows various corrosion phenomena observed inside a corroded pin. The black regions in the grey aluminium matrix are corroded areas, whereas the small white spots represent intermetallic particles. Fig. 5a shows a 3D reconstruction including LT and L planes. Two types of attack can be observed; one is a typical intergranular corrosion (IGC), while the other is referred to as an ‘‘exfoliation-like attack” (ELA). The term ‘‘exfoliation-like” corrosion is used to indicate morphologies similar to exfoliation corrosion. In contrast to conventional exfoliation corrosion, ELA is not present along the surface and the mechanism might be different. IGC, seen as thin black paths within the Al matrix, follows the grain boundaries without indication of a specific preferred direction as would be expected for elongated grains [34]. Here the grain boundaries are selectively attacked and single grains may be completely surrounded by IGC, but single corrosion paths of several 100 lm in length along grain boundaries can also be observed. ELA-type corrosion can start in the bulk of the material. Fig. 5a illustrates this and the high aspect ratio (length to width) of the ELA. It is also clear that ELA is strongly directed and follows the L direction in planes parallel to the former sheet surfaces (orthogonal to ST direction). A preferred location for ELA at the pin surface or pin center was not found. A further observation is that ELA is not hindered by grain boundaries and can be in the range of mm length (Fig. 5c), while its thickness is in the range of 10 lm. In Fig. 5b, a top view inside the pin supports the observations made in Fig. 5a and shows the differing behaviour of IGC and ELA clearly. The restriction of ELA to about 10 lm width is obvious. The long IGC paths follow random directions. Intersections of IGC and ELA can be observed. At such locations the direction of the ELA is not disturbed, but nevertheless some ELA paths coincide with IGC. Fig. 5c shows an ST plane view (plane parallel to the sheet surface) on the corroded pin. Again, IGC and ELA are visible. It is important to note that the surface of the pin is heavily subjected to localized attack but still has enough mechanical integrity to remain in place. The matrix below this surface layer is heavily dissolved. As the machined layer was not removed from the pin after production, this undissolved surface layer can be attributed to the mechanical deformed layer.
Fig. 5. Reconstructed X-ray tomograms. Various ex situ views of corrosion sites within an AA6111 pin exposed to 0.7 M HCl solution for 7 h: (a) 3D view and section; (b) L plane (top view) taken out of the lower region of the pin; and (c) ST plane (parallel to sheet surface) showing the whole pin. Clearly visible is the exfoliation-like attack, the intergranular corrosion and the special behaviour of the surface–deformed layer.
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Fig. 6 presents L planes (top view) from two different pin heights (top and bottom) of alloys AA6111 ( Fig. 6a and b) and AA6016 ( Fig. 6c and d). The morphology of corrosion attack changes along the height of the pin. In the upper part IGC is more pronounced, while approaching the bottom part ELA outweighs IGC. Comparing the two alloys, it is seen that alloy AA6111 ( Fig. 6a and b) is more susceptible to ELA attack than AA6016 (Fig. 6c and d). In general it is striking that the grain morphology and size revealed by the IGC attack has no relation to the ELA path distances or directions. While the grains are not elongated, the ELA has a very high aspect ratio and propagates directly through grains and the grain boundaries. Therefore it is assumed that ELA propagation is not controlled by grain boundaries. The images of alloy AA6016 (Fig. 6c and d), however, indicate that the ELA can initiate and develop out of an IGC path. As exemplary illustrated inFig. 6c and d the ELA paths are not in direct contact with the surface. They are only in contact with the solution present in the IGC path. 3.2. Corrosion in 0.7 M HCl solution characterized in situ In this section, the ability of the method to monitor corrosion processes in situ is demonstrated and used to analyze, in three dimensions, intergranular corrosion propagation rates during immersion of the pin. The obtained data are compared to measurements obtained by conventional 2D analysis.
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Fig. 7 presents in situ X-ray microtomography sections for an AA6111 pin immersed in 0.7 M HCl solution. The corrosion propagation is monitored by X-ray tomography during the whole immersion period of 6 h. The open circuit immersion corrosion experiment was not halted during the tomography measurement, and the propagation rates of the corrosion can be estimated from the sequence of images. After an initiation time of 2–3 h during which no attack can be detected the corrosion starts at the surface. After 4 h of immersion significant IGC propagation can be seen and additional initiation sites appear. With increasing immersion time the IGC corroded volume increases. (Note: compared to ex situ experiments the overall corrosion volume was less pronounced in in situ experiments most probably due to experimental conditions, in particular bubble formation on the surface). Fig. 8 presents the corroded volume obtained from the experiments illustrated in Fig. 7. To measure corroded volumes, the tomograms were converted into binary images and a greyscale threshold was defined to allow separation of corroded and not-attacked regions. It can be seen that after an initiation time of >2 h in-depth corrosion propagation occurs. After initiation the increase of the corrosion volume is constant, indicating an approximately constant IGC growth rate. The two curves in Fig. 8 additionally show the difference in the corrosion volume estimated on the one hand by a labour-extensive cross-sectioning experiment (2D), and on the other hand measured using 3D tomography.
Fig. 6. Morphology of attack characterized ex situ after immersion in 0.7 M HCl for 7 h depending on the position within the pin for AA6111: (a) upper section and (b) lower section and AA6016, (c) upper section and (d) lower section.
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Fig. 7. In situ microtomograms of AA6111 immersed in 0.7 M HCl. Cross-sections from the middle position of the pin showing the corrosion morphology development as a function of immersion time, 2 h (a), 4 h (b) and 6 h (c).
Five spatially separated cross-sections of the tomogram (reconstructed image of the pin) were examined to simulate the 2D crosssectioning experiment that would be obtained by SEM observation of metallographic sections. For the 3D analysis of the corroded volume, the whole pin was considered. The measured corrosion volumes were similar for both estimations. Nevertheless, the 3D method detects more corrosion in general because small fractions of materials with a high percentage of corroded volume are better resolved. The ability to detect non-uniformly distributed corrosion in samples is a major advantage of the X-ray tomography method compared to conventional 2D methods. 3.3. Corrosion in 2.5 M HCl solution characterized in situ
Fig. 8. AA6111 exposed to 0.7 M HCl for 7 h in situ monitored. Comparison of the corroded volume estimated by the analysis of five cross-sectional slices (2D) and measured in the volume measured in 3D.
To find an explanation for the transition from intergranular corrosion (IGC) to exfoliation-like attack (ELA) and the occurrence of ELA in general, experiments were performed in a more aggressive environment. Based on the assumption that in IGC paths more aggressive solutions build up and based on the observation that ELA initiates at IGC sites, ELA seems to be more pronounced in aggressive solutions. A pin of each alloy was immersed in 2.5 M HCl. The corrosion process was then monitored in situ using a different tomography setting that enables faster data acquisition.
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These settings gave a reasonable balance between signal-to-noise ratio and acquisition time (10 min per scan). Fig. 9 presents a cross-section from the middle of an AA6111 pin immersed in 2.5 M HCl. After an initiation time of about 30 min, rapid localized corrosion showing mainly the directed ELA morphology occurs, but IGC is also visible. IGC apparently did not propagate further between 45 min and 1 h of solution exposure. ELA, however, grew within the ST plane and new ELA sites occurred within this time span. The broadening of the ELA corrosion path is initially fast but stops at a final thickness of about 10 lm. Fig. 10 presents a comparison of the corrosion morphology observed for alloys AA6016 and AA6111 during immersion experiment in 2.5 M HCl. Both alloys show mainly ELA, but as already reported for the less aggressive solution, AA6016 alloy is significantly less susceptible to corrosion. This is derived from a longer initiation time (AA6016 1 h; AA6111 30 min) and the extent of corrosion after 2 h of immersion. Again ELA showed no significant corrosion path broadening for longer immersion times and restricted propagation in ST planes (parallel to sheet surface). These observations imply that like the IGC, the attack may follow a certain microstructural feature. Possible microstructural features that may explain this phenomenon are: (i) grain boundaries and texture, (ii) alignment of the intermetallics and (iii) composition variations in the alloy. In the following of these topics only the influence of grain boundaries and texture
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on ELA is further analysed while a detailed discussion of the ELA mechanism is described in a different paper [42]. 3.4. Exfoliation-like attack (ELA) and grain boundaries/texture Fig. 11a presents a typical EBSD map of a corroded AA6111 (0.7 M HCl, 7 h, ex situ) pin. The different colours represent various grain orientations. The grain boundaries are orientated slightly along the LT direction and no preferential texture is present in the alloy. Fig. 11b shows a SEM image of the identical section. EBSD map in Fig. 11a, revealing most of the grain boundaries, is overlaid on the SEM image. Two sites of ELA, IGC (black) and intermetallics (bright spots) are visible. The IGC follows the grain boundaries, whereas the ELA orientation does not correspond to the grain boundaries. Fig. 12a presents another experimental approach proving the independence of ELA from the grain boundaries in sheet material. It shows a SEM image of a surface of alloy AA6016 exposed to 5 M HCl for 15 min (the increased aggressivity of the solution used for this surface attack followed from the observation that for very acidic conditions the directed ELA strongly dominates the IGC attack). The image was taken from the L plane, looking onto the sheet thickness. It can be seen that the ELA is very distinct and has spacing between the ‘‘lamella” of about 10 lm, which corresponds to
Fig. 9. AA6111 cross-section from the middle position of the pin showing corrosion development as a function of immersion time in 2.5 M HCl solution. IGC is present, and ELA is dominant. One area is magnified to show the IGC and exfoliation-like corrosion together.
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Fig. 10. Comparison of attack in AA6016 and AA6111 after exposure to 2.5 M HCl solution. For both alloys the morphology is similar but the time scale differs.
the ELA path thickness reported previously. The lamella periodicity is significantly smaller than the grain size, and therefore this attack morphology cannot be provoked by the grains or, respectively, the grain boundaries. Fig. 12b shows an EBSD map of the corroded sheet shown in Fig. 12a after polishing away most of the corrosion ‘‘lamella”. The corrosion sites (ELA) visible in black do not follow the grain boundaries. Hence in this experiment, too, ELA did not depend on texture or grain boundaries. 3.5. Dissolution of intermetallic particles within the material
Fig. 11. AA6111 pin corroded in 0.7 M HCl for 7 h (ex situ). (a) EBSD mapping AA6111 showing the grain boundaries and orientation. The pin was polished after corrosion to perform EBSD. (b) SEM/EBSD image; intergranular corrosion (IGC) and exfoliation-like attack (ELA) is visible. Grain boundaries (GB) are marked in white. ELA follows the ST planes.
The behaviour of intermetallics in an aggressive solution was analysed by X-ray microtomography, revealing a new phenomenon. A complete dissolution of micrometer-sized intermetallics, which are reported to be cathodic when characterized at the surface, was observed [43,44]. Fig. 13 illustrates this corrosion attack inside the bulk alloy. The white microscopic features seen within the grey dark matrix are areas of higher absorption coefficient related to the presence of intermetallics. For the alloy presented, these bright features are Al5FeSi and Al15(Fe,Mn)3Si2 intermetallics which also contain a non-stoichiometric amount of copper. Fig. 14 shows the predicted equilibrium composition and intermetallics for the
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4. Discussion Different types of in-depth corrosion propagation phenomena have been successfully investigated by means of X-ray microtomography. The following section discusses different aspects of these types of attack in relation to literature data. 4.1. Surface–deformed layer
Fig. 12. AA6016;exposed to 5 M HCl solution for 15 min: (a) A SEM image of a heavily-corroded surface facing the longitudinal and transverse direction (LT plane). Small aligned lamella (ELA) are visible. (b) EBSD mapping after polishing away most of the lamella (ELA). It can be seen that the exfoliation-like corrosion path is not oriented along the grain boundaries.
alloys used. The thermodynamic calculation model is based on the PANDAT database [45,46]. The black vertical bar shows the final ageing treatment temperature of 180 °C (for 45 min), from which the alloy is quenched to room temperature. For AA6016 (Fig. 14a) mainly Mg2Si, Si and Q(Al5Cu2Mg8Si6) phase will form nanometer-sized [25,47]. At these low homologous temperatures, diffusion is slow and therefore no equilibrium will be reached in 45 min. In addition, primary AlFeSi intermetallics with an average size of about d = 3.5 lm (80% < 4 lm) are present directly from the melt and represent a surface coverage of about 0.8% for these alloys. AA6111 contains primary intermetallics similar in shape and size to those in AA6016 but with a volume content of 1.4%. The main difference from the point of view of corrosion for AA6111 is the non-stoichiometric Cu content in the primary AlFeSi intermetallics, as well as the lower Si content in solid solution. The predicted composition was experimentally verified with electron microprobe (EPMA) measurements. The pin displayed in Fig. 13 was exposed to 0.7 M HCl for 5 h. During this exposure period the Fe containing intermetallics dissolve when an intergranular corrosion path and therefore aggressive solution enriched with different dissolved metallic elements (i.e. MeClx,H+) reaches the intermetallics. After 1 h of exposure, all the intermetallic particles highlighted are still present and no sign of corrosion is detected. After 3 h of immersion, the first corrosion paths intersect with intermetallics. Those intermetallics in contact with the solution leave a black area, indicating that the intermetallics have dissolved. After 5 h of immersion, heavy traces of corrosion at the former intermetallic sites are visible. Hence it may be concluded that cathodic Fe containing intermetallics can dissolve in the inner part of materials during the corrosion process and that local aggressivity can be very high in IGC attack.
At an initial stage, a surface–deformed layer resulting from the machining process is more susceptible to localized corrosion compared to the bulk structure, as reported in literature for submicrometer thick surface layers [27,30,48]. This susceptibility is attributed to the higher defect density present at the surface (e.g. impurities, oxide inclusions, contamination and ultrafine grain structure) [27]. The influence of this deformed layer on filiform corrosion under an organic coating is also discussed in literature [27,49], where it is shown that this type of layer is attacked first and causes further delamination of the coating for certain heat treatments. But a distinction must be made in this case between superficial filiform corrosion and deep penetrating attack. This study was able to show that after the first preferential IGC attack the deformed layer surface does not dissolve further. This indicates a high uniform corrosion resistance in this very aggressive electrolyte. Nevertheless, because this layer is very thin, it is fragile and difficult to detect using conventional methods. This observation raises a question concerning the positive influence of plastic deformation or of the decrease in grain size on corrosion resistance especially for in-depth corrosion propagation. 4.2. ELA vs. IGC and correlation of ELA with grain boundaries To summarize the corrosion initiation and propagation experiments, two types of attack were observed: the IGC and the exfoliation-like attack (ELA). The initiation time is comparable for IGC and ELA whereas ELA develops faster. The results of the experiments performed in the two solutions (0.7 M and 2.5 M HCl) facilitate discussion of the two corrosion phenomena. The assumption that ELA requires higher aggressivity is verified; the apparent aggressivity threshold lies between the two solutions chosen for the experiments. In 2.5 M HCl the two corrosion types compete, but ELA penetrates more deeply (and therefore more quickly) than IGC. In addition, susceptibility of alloy AA6111 to ELA and IGC is higher than that of AA6016. The main differences between AA6016 and AA6111 are the small grain size structure of AA6016 in the center of the pin, higher excess of Si of AA6016 and higher Cu/Mg content of AA6111, which is reflected in a non-stoichiometric Cu content in the intermetallics as well as Cu in solid solution for AA6111. The grain structure is not further considered in this paper as ELA is not seen to depend on the grain boundaries. In AA6016, fine Si-containing intermetallics are present and Si is present in larger amounts in solid solution. Taking these microstructural features into account in relation to the difference between the corrosion susceptibilities of AA6111 and AA6016, it may be concluded that the intermetallics and/or the solid solution content play an important role in the ELA process. The grain boundaries are not responsible for the direction of ELA, especially in aggressive solutions. In a first comparison of intermetallics alignment and ELA corrosion path with the help of X-ray microtomography, no significant correlation was found. Independent of the aggressivity of the bulk solution in all experiments ELA was directed parallel to the former sheet surface (ST plane) and showed a final thickness of 10 lm. These observations are a strong indication that ELA might be correlated to compositional variations in the
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Fig. 13. Sequence of tomograms obtained for AA6111 after 1 h, 3 h and 5 h exposure to 0.7 M HCl. ‘‘Cathodic” intermetallics (white spots) dissolve. After 5 h exposure almost all intermetallics in contact with solution dissolved.
matrix (i.e. Mg, Si and Mn). It is matter of a detailed investigation [42]. 4.3. Dissolution of intermetallics In corrosion experiments where surface processes are considered, the fact that Fe containing intermetallics disappear is mainly attributed to undercutting as a result of dissolution of the surrounding matrix [50], the so called ‘‘trenching process”. It is obvious from Fig. 13 that the corrosion mechanism is different within the bulk of the material, where various solution concentrations induced by the confinement of corrosion products can be found. Trenching around an intermetallic may still be envisaged, but has not been observed in the X-ray microtomography experiments. The complete dissolution of the intermetallics seen can be explained by the fact that if they are not in direct electrical contact with the surrounding Al matrix, they are no longer cathodically protected and can freely corrode. This dissolution can also occur as a result of chemical dissolution due to the extreme aggressivity of the local solution. What has to be stated here is that this is not only a dealloying process often observed for intermetallic particles on the surface and discussed in the introduction, but a complete dissolution. In the case of dealloying, the contrast in microtomography would subsist as it is not the most noble and heavy element that would dissolve. As a result, Fe ions as well as Cu ions are expected to be present in the solu-
tion and can redeposit as further cathodic sites at different locations or act as strong oxidation agents and accelerate the corrosion process. In literature so far it has been stated that the Fe containing intermetallics at the surface undergo selective dissolution of Al, leaving a reactive Fe-, Mn- and Si-enriched remnant. While Fe is said to provoke a higher cathodic activity, Mn and Si are supposed to reduce the cathodic activity because of the formation of an ‘‘inert” passive layer [44]. For AA2024 alloy, extensive work was also performed on Cu-containing S phase particles, where Mg and Al dissolve leaving a free-corroding Cu sponge-like structure with subsequent dissolving Cu redepositing on the surrounding Al matrix [43,51]. Both processes generate greater cathodic activity on the part of the intermetallics in the corrosion process. All the phenomena described do not apply for the complete dissolution of the intermetallics observed, or at least cannot be resolved time-wise if the dissolution is the result of multiple dissolution steps. X-ray microtomography, however, offers an excellent opportunity to detect the complete dissolution of intermetallics within the bulk as a function of time, which is extremely hard or impossible to verify with conventional methods. The location of dissolved intermetallics is detected due to the 3D reconstruction, and the corrosion process can be followed as a function of time within the bulk without taking into account artefacts produced by sectioning or sample preparation. These results raise the question of the chemistry within an intergranular corrosion path.
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path is highly aggressive. It also indicates that the corrosion mechanisms within the bulk may differ from the corrosion attack observed and well-documented at the surface. Acknowledgements The authors would like to thank P. Wyss, Empa Dubendorf, for support in image processing. They are indebted to Axel Stange, Alcan Inc., and Tom Quested, Novelis AG, for help with EBSD measurements and sample preparation. Financial support from the Alcan Technology Fund is gratefully acknowledged. The use of the Advanced Photon Source was supported by the US Department of Energy, Office of Science, Office of Basic Energy Sciences, under Contract No. DE-AC02-06CH11357. References
Fig. 14. Thermodynamic modelling of the equilibrium composition of the alloys with regard to their intermetallic phases for (a) AA6016 and (b) AA6111. The dotted black bar marks the first heat treatment temperature and the black bar at 180 °C marks the final heat treatment temperature.
5. Conclusion In this study the corrosion processes of two AlMgSi alloys were monitored in three dimensions. The synchrotron X-ray microtomography technique was used for in situ characterization. Different issues concerning corrosion morphology and propagation rates were documented, and are discussed: 1. X-ray l-tomography is a useful tool to monitor in situ corrosion processes in the lm-range. Due to the unique possibility of non-destructively viewing inside a large volume of material in <10 min for each scan, time series of corrosion processes characterization within the bulk can be performed and displayed in three dimensions. This facilitates the description of phenomena which are almost impossible to detect directly using conventional methods. 2. Surface–deformed layers are seen to be very susceptible to intergranular corrosion, but after first attack occurs, they are not uniformly corroded further, even though the matrix underneath dissolves completely. 3. A new type of exfoliation-like attack (ELA) was observed experimental with the bulk of the material. The actual mechanism is not yet clear, but the attack does not follow the grain structure, texture or planes of primary intermetallics. ELA was more pronounced for fast corrosion processes in aggressive solution, while for slow processes in less aggressive solutions intergranular corrosion (IGC) prevailed. 4. Intermetallics supposed to be cathodic are seen to dissolve during the corrosion process within the bulk material. This observation provides a hint that the solution within a corrosion
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Contents lists available at ScienceDirect
Corrosion Science journal homepage: www.elsevier.com/locate/corsci
In situ monitoring of corrosion processes within the bulk of AlMgSi alloys using X-ray microtomography Fabian Eckermann a,b,*, Thomas Suter a,*, Peter J. Uggowitzer b, Andreas Afseth c, Alison J. Davenport d, Brian J. Connolly d, Magnus Hurlen Larsen e, Francesco De Carlo f, Patrik Schmutz a a
Laboratory for Corrosion and Materials Integrity, Empa, Swiss Federal Laboratories for Materials Testing and Research, Uberlandstrasse 129, 8600 Dubendorf, Switzerland Laboratory of Metal Physics and Technology, Department of Materials, ETH Zurich, 8093 Zurich, Switzerland c Alcan Inc., Voreppe 38341, France d University of Birmingham, Metallurgy and Materials, Edgbaston, Birmingham, B15 2TT, UK e Department of Materials Science and Engineering, Norwegian University of Science and Technology, N-7491 Trondheim, Norway f Advanced Photon Source, Argonne National Laboratory, Argonne IL, United States b
a r t i c l e
i n f o
Article history: Received 29 April 2008 Accepted 2 September 2008 Available online 23 September 2008 Keywords: A. Aluminium B. X-ray microtomography C. Acid corrosion C. Exfoliation corrosion
a b s t r a c t Susceptibility to localized corrosion of AlMgSi (AA6016 and AA6111) alloys in certain aggressive environments is high. In this study, synchrotron X-ray microtomography was used to monitor non-destructively corrosion processes within bulk materials. In the selected aggressive solutions, surface–deformed layers showed high localized corrosion susceptibility, but are more stable than the bulk of the alloy during corrosion propagation. In addition, exfoliation-like attack was observed as a transition from intergranular attack. This directed corrosion is not related to grain or crystallographic structure. Finally, intermetallic particles dissolution inside the materials after contact with aggressive solution of the intergranular corrosion path was evidenced. Ó 2008 Elsevier Ltd. All rights reserved.
1. Introduction X-ray microtomography is a non-invasive technique used to image three-dimensional (3D) bulk objects. Visualization of defects, dimensional inspection and local characterization are possible in the interior of a material [1]. The spatial resolution of the method can vary from several mm to the 100 nm scale. Use of higher spatial resolution limits the observed volume and acquisition times [2–7]. Measurement times also range from several hours to a few seconds depending on the source, detector, material and specimen dimensions. Over the past five to ten years, X-ray microtomography has become an increasingly accepted technique for qualitative and quantitative characterization of the 3D internal structures of materials. A very new field of tomography research using this capability is the study of corrosion processes. Whereas an initial in situ tomography experiment was performed in 2004 involving the corrosion of sandstone under triaxial compression and chemical atmosphere
* Corresponding authors. Address: Laboratory for Corrosion and Materials Integrity, Empa, Swiss Federal Laboratories for Materials Testing and Research, Uberlandstrasse 129, 8600 Dubendorf, Switzerland (F. Eckermann). Tel.: +41 44 823 4921; fax: +41 44 823 4015 (T. Suter). E-mail addresses: [email protected], [email protected] (F. Eckermann), [email protected] (T. Suter). 0010-938X/$ - see front matter Ó 2008 Elsevier Ltd. All rights reserved. doi:10.1016/j.corsci.2008.09.015
with a resolution of 0.1225 mm3 per voxel (=3D pixel) [8], studies of corrosion processes of metals have only started recently [9–11]. In corrosion studies it is of crucial importance to understand the changes occurring within the bulk of a material when it is exposed to an aggressive aqueous medium. In many localized corrosion sites of attack most damage takes place beneath the surface. However, most models describing localized corrosion have been deduced from ‘‘electro”chemical activity measured on the surface of a material or via traditional 2D, ex situ destructive analysis of metallographic cross-sections. From conventional methods it is difficult to obtain reliable information on complex corrosion propagation morphologies and on the interaction of microstructure with localized corrosion or corrosion cracking. The use of conventional methods (e.g. polishing) also enhances the risk of introducing artefacts during sample preparation. Grain fallout during polishing is one specific example. In the first publication concerning in situ characterization of corrosion for a metal, Marrow et al. [9] observed stress corrosion cracking development in stainless steel. Corrosion processes could be quantified and a 3D model was established. Recently Davenport et al. [10] used synchrotron smonochromatic X-rays (energy resolution of DE/E of 1.4 10 4) to correlate the yttrium distribution in a magnesium alloy with the corrosion path. In the first published work on in situ aluminium corrosion investigated using X-ray microtomography, Connolly et al. [11] demonstrated some of
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the corrosion phenomena occurring on an AA2024 alloy, such as intergranular corrosion path propagation and the dissolution of the matrix around intermetallics within AA2024. In this work three phenomena related to the corrosion of Al– Mg–Si alloys are investigated using X-ray microtomography: The role of intermetallic particles, the behaviour of a surface–deformed layer and the occurrence of exfoliation-like attack (ELA). 1.1. The role of the primary intermetallics in the corrosion process Considerable work has been carried out characterizing the cathodic trenching of intermetallic particles [12–14], often following initial dealloying reactions [15–21]. The role of certain intermetallic particles in pit initiation is also well established [17,18,22–26]. However, very few observations have been made of the behaviour of particles in the interior of materials, where aggressive solutions can build up in localized corrosion sites. 1.2. The behaviour of the surface–deformed and activated layer A surface–deformed layer is produced during any shearing event which takes place on the surface of aluminium, i.e. grinding, rolling or machining. This surface layer can induce filiform corrosion susceptibility depending on the formation of second phase particles and a fine grain structure [27]. The second phase particles are formed during heat treatments after the deformation process [27–30]. In the cited work this layer is of submicrometer size. However concepts presented for this surface–deformed layer might be transferred to surface–deformed layers of much larger thickness as present in this study. The stability of the surface–deformed layer after initiation of first localized attack has not so far been addressed in detail in literature. 1.3. Exfoliation-like attack (ELA) Exfoliation is a corrosion process assumed to occur in highly directional microstructures. A consensus on its mechanism has not yet been reached, but it has been attributed variously to intermetallics arranged in planes, to mechanisms similar to those for intergranular corrosion and stress corrosion cracking (in the 2xxx series), and also to segregation of alloying elements coming from the ingot [31,32]. Frankel stated for the AA2024 and AA7xxx alloys using a combination of different characterization techniques such as X-ray radiography, foil penetration technique and potentiodynamic polarization measurements that exfoliation follows intergranular corrosion and that the intergranular corrosion process moves along elongated grains faster in certain directions [33]. The kinetics of intergranular corrosion also depends on the grain aspect ratio [34]. Kelly claimed that exfoliation is triggered by stresses generated by corrosion products and follows an intergranular corrosion path [35]. Others have attributed exfoliation corrosion to segregation processes. The directed microstructure in the final rolled sheet is, for example, a remnant of the ingot microstructure [31,36,37]. None of the above-mentioned corrosion phenomena have so far been monitored and discussed for AlMgSi alloys in the context of X-ray microtomography investigations. This method has the advantage of monitoring of corrosion attack and tracking of corrosion paths as a function of time during corrosion processes which occur upon aggressive solution contact. It also makes it possible to characterize the morphology of corrosion propagation. In this way, not-yet-clarified changes in the behaviour of surface–deformed layers and intermetallics during corrosion propagation can be investigated. New information improving the understanding of the role of intermetallics, surface–deformed layers, corro-
sion morphology and indirectly about the solution chemistry can be gained. 2. Experimental X-ray microtomography experiments were performed at the Argonne National Laboratories, IL, USA, Advanced Photon Source beam line 2-BM [38]. Acquisition was carried out at a 17 keV Xray energy with a double Si 1 1 1 crystal monochromator. Fast acquisition of the radiographs was performed at 20 keV with a double multilayer monochromator. The detector CCD chip was read out with 1024 1024 pixels, giving a final practical resolution of about 3 lm. Seven hundred and twenty 2D radiographs were taken for the subsequent reconstruction procedure. The settings used for fast acquisition of the data lowered the acquisition time from 1 h to 10 min. For image processing, ImageJ software [39,40] was used. To measure corroded volumes, the tomograms were converted into binary images and a greyscale threshold was defined to allow separation of corroded and not-attacked regions. The set-up is designed so that a pin can be fully immersed in solution during tomographic measurements, whereby the dimensions of the pin are adjusted to allow short acquisition times (see Fig. 1). Samples for X-ray tomography experiments were first machined to 3 mm diameter rods, and then the upper 2 mm was turned down to 0.5 mm diameter pins. This thin part was analysed with X-ray tomography and is positioned in the middle of the sheet. The samples were lacquered with Stopping Off Lacquer to cover all but the to the X-ray beam exposed part of the pin. For samples investigated in situ, a silicone rubber tube with 3 mm internal diameter was slipped over the outside of the 3 mm rod to form a cell around the pin, and the cell was subsequently filled with solution. For the electron backscattered diffraction measurements (EBSD), a LEO GEMINI 1530 field emission gun SEM was used. The data were acquired at an angle of 70° with respect to the surface and 15 kV electron beam voltage. The aperture of the electron beam optics was opened to 120 lm. For corrosion tests performed on the sheet material, the surface was ground with 4000 SiC paper before being exposed to the solution.
Fig. 1. Sketch of exposure cell: a pin of 500 lm diameter was fully immersed in the solution confined in silicone rubber tubing. The field of view marks the area where the beam penetrated the specimen.
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2.1. Materials: AA6016 and AA6111 alloys The composition of aluminium alloys 6016 and 6111 is given in Table 1. The pins for the X-ray tomography experiments were cut from the center of 3 mm-thick aluminium sheets. The machined layer (surface–deformed) was not removed prior to the corrosion experiments. All samples used in this study were heat treated to 540 °C for 15 min (solution heat treatment) after machining and then water-quenched. Subsequent ageing treatment took place for 45 min at 180 °C, followed by water-quenching to achieve high intergranular corrosion susceptibility. For the EBSD mappings performed after corrosion experiments, the pin surfaces were polished to 0.05 lm with a vibromed polishing system. Fig. 2 shows the microstructure of the aluminium alloys used. Considering the cross-section through the AA6016 sheet (Fig. 2a), the microstructure obtained was very heterogeneous. Large grains in mm size dominate in the outer region; the center of the sheet consists of very small grains in the 10 lm-range. A similar heterogeneity was observed by Svenningsen et al. [41] for extruded AlMgSi, where large grains of over 1mm in size dominated the surface of the sheet, while in the center region fine grains were present. The nomenclature used for labelling the orientations is visualized in Fig. 3. No significant difference between the longitudinal (L) plane and longitudinal–transverse (LT) plane can be found in the grain structure. Within the AA6016 pins, the mentioned heterogeneous grain structure (fine and coarse grains) was present because the pins were cut out of the middle of the sheet material (shown in Fig. 3). In alloy AA6111 a different grain structure can be identified (Fig. 2b). The grains are slightly directed (aspect ratio of 3.6), and average a thickness of about 120 lm. Grain structure did not differ in size and shape in the LT or L plane.
Fig. 3. Schematic sketch showing the position of the pin in the sheet and the nomenclature used.
To investigate the corrosion processes of the AA6016 and AA6111 alloys, three types of corrosion experiments in particular were performed: (1) in situ with samples immersed in 0.7 M HCl; (2) in situ in 2.5 M HCl; and (3) ex situ after 0.7 M, 2.5 M and 5 M HCl solution exposure. The immersion times and solutions are listed in Table 2. AA6111 was immersed in 0.7 M HCl solution for up to 7 h and monitored in situ with 1 scan (X-ray tomography measurement) per hour. In situ experiments with specimens immersed in 2.5 M HCl were monitored with 1 scan every 15 min. Ex situ experiments were investigated after exposure to solution. The solution was refreshed after each scan to minimize unwanted effects as a change in the solution pH. Table 2 Overview of the experiments presented in this paper Experiment type
AA6016 (immersion time [h])
AA6111 (immersion time [h])
In situ
–
7
2
1
0.75 7 45 0.25 2 0.25
– 7 10 – 1 –
Ex situ Table 1 Composition of AA6016 and AA6111 measured by optical emission spectroscopy Wt%
Mg
Si
Fe
Mn
Cu
Cr
AA6016 AA6111
0.35 0.61
1.05 0.80
0.19 0.26
0.08 0.21
0.07 0.70
0.01 0.02
Ex situ Ex situ
0.7 M HCl 2.5 M HCl 0.7 M HCl 2.5 M HCl 5 M HCl
Type of material, immersion time and solution type are listed.
Fig. 2. Optical image of the alloys’ microstructure after Barker’s etching: (a) For AA6016 sheet material a significant difference in the grain structure through the sheet thickness is observed. (b) For AA6111 oriented the same way, homogeneous grain size distribution is visible.
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3. Results 3.1. Overview of the observed attack In this section, the newly-revealed phenomena of slow-dissolving surface–deformed layer and the transition from intergranular corrosion (IGC) to exfoliation-like attack (ELA) are illustrated. In Fig. 4, an ex situ SEM image of a corroded AA6016 pin used in the tomography experiments is shown, providing an overview of the corrosion damage observed upon immersion in aggressive HCl electrolyte. The pin (Fig. 4) was immersed for 45 h in 0.7 M HCl. The previously mentioned heterogeneous grain structure can be seen, as the pin is not attacked homogeneously. The preferred re-
Fig. 4. SEM image of a pin with heavy traces of corrosion. AA6016 immersed for 45 h in 0.7 M HCl.
gion of attack in the middle of the pin corresponds to the smallgrained section (Fig. 2a). However, the extent and morphology of the corrosion attack inside the pin is not predictable from such surface observation. In X-ray microtomography experiments the corrosion attack morphology within pins is visualized. Fig. 5 shows various corrosion phenomena observed inside a corroded pin. The black regions in the grey aluminium matrix are corroded areas, whereas the small white spots represent intermetallic particles. Fig. 5a shows a 3D reconstruction including LT and L planes. Two types of attack can be observed; one is a typical intergranular corrosion (IGC), while the other is referred to as an ‘‘exfoliation-like attack” (ELA). The term ‘‘exfoliation-like” corrosion is used to indicate morphologies similar to exfoliation corrosion. In contrast to conventional exfoliation corrosion, ELA is not present along the surface and the mechanism might be different. IGC, seen as thin black paths within the Al matrix, follows the grain boundaries without indication of a specific preferred direction as would be expected for elongated grains [34]. Here the grain boundaries are selectively attacked and single grains may be completely surrounded by IGC, but single corrosion paths of several 100 lm in length along grain boundaries can also be observed. ELA-type corrosion can start in the bulk of the material. Fig. 5a illustrates this and the high aspect ratio (length to width) of the ELA. It is also clear that ELA is strongly directed and follows the L direction in planes parallel to the former sheet surfaces (orthogonal to ST direction). A preferred location for ELA at the pin surface or pin center was not found. A further observation is that ELA is not hindered by grain boundaries and can be in the range of mm length (Fig. 5c), while its thickness is in the range of 10 lm. In Fig. 5b, a top view inside the pin supports the observations made in Fig. 5a and shows the differing behaviour of IGC and ELA clearly. The restriction of ELA to about 10 lm width is obvious. The long IGC paths follow random directions. Intersections of IGC and ELA can be observed. At such locations the direction of the ELA is not disturbed, but nevertheless some ELA paths coincide with IGC. Fig. 5c shows an ST plane view (plane parallel to the sheet surface) on the corroded pin. Again, IGC and ELA are visible. It is important to note that the surface of the pin is heavily subjected to localized attack but still has enough mechanical integrity to remain in place. The matrix below this surface layer is heavily dissolved. As the machined layer was not removed from the pin after production, this undissolved surface layer can be attributed to the mechanical deformed layer.
Fig. 5. Reconstructed X-ray tomograms. Various ex situ views of corrosion sites within an AA6111 pin exposed to 0.7 M HCl solution for 7 h: (a) 3D view and section; (b) L plane (top view) taken out of the lower region of the pin; and (c) ST plane (parallel to sheet surface) showing the whole pin. Clearly visible is the exfoliation-like attack, the intergranular corrosion and the special behaviour of the surface–deformed layer.
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Fig. 6 presents L planes (top view) from two different pin heights (top and bottom) of alloys AA6111 ( Fig. 6a and b) and AA6016 ( Fig. 6c and d). The morphology of corrosion attack changes along the height of the pin. In the upper part IGC is more pronounced, while approaching the bottom part ELA outweighs IGC. Comparing the two alloys, it is seen that alloy AA6111 ( Fig. 6a and b) is more susceptible to ELA attack than AA6016 (Fig. 6c and d). In general it is striking that the grain morphology and size revealed by the IGC attack has no relation to the ELA path distances or directions. While the grains are not elongated, the ELA has a very high aspect ratio and propagates directly through grains and the grain boundaries. Therefore it is assumed that ELA propagation is not controlled by grain boundaries. The images of alloy AA6016 (Fig. 6c and d), however, indicate that the ELA can initiate and develop out of an IGC path. As exemplary illustrated inFig. 6c and d the ELA paths are not in direct contact with the surface. They are only in contact with the solution present in the IGC path. 3.2. Corrosion in 0.7 M HCl solution characterized in situ In this section, the ability of the method to monitor corrosion processes in situ is demonstrated and used to analyze, in three dimensions, intergranular corrosion propagation rates during immersion of the pin. The obtained data are compared to measurements obtained by conventional 2D analysis.
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Fig. 7 presents in situ X-ray microtomography sections for an AA6111 pin immersed in 0.7 M HCl solution. The corrosion propagation is monitored by X-ray tomography during the whole immersion period of 6 h. The open circuit immersion corrosion experiment was not halted during the tomography measurement, and the propagation rates of the corrosion can be estimated from the sequence of images. After an initiation time of 2–3 h during which no attack can be detected the corrosion starts at the surface. After 4 h of immersion significant IGC propagation can be seen and additional initiation sites appear. With increasing immersion time the IGC corroded volume increases. (Note: compared to ex situ experiments the overall corrosion volume was less pronounced in in situ experiments most probably due to experimental conditions, in particular bubble formation on the surface). Fig. 8 presents the corroded volume obtained from the experiments illustrated in Fig. 7. To measure corroded volumes, the tomograms were converted into binary images and a greyscale threshold was defined to allow separation of corroded and not-attacked regions. It can be seen that after an initiation time of >2 h in-depth corrosion propagation occurs. After initiation the increase of the corrosion volume is constant, indicating an approximately constant IGC growth rate. The two curves in Fig. 8 additionally show the difference in the corrosion volume estimated on the one hand by a labour-extensive cross-sectioning experiment (2D), and on the other hand measured using 3D tomography.
Fig. 6. Morphology of attack characterized ex situ after immersion in 0.7 M HCl for 7 h depending on the position within the pin for AA6111: (a) upper section and (b) lower section and AA6016, (c) upper section and (d) lower section.
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Fig. 7. In situ microtomograms of AA6111 immersed in 0.7 M HCl. Cross-sections from the middle position of the pin showing the corrosion morphology development as a function of immersion time, 2 h (a), 4 h (b) and 6 h (c).
Five spatially separated cross-sections of the tomogram (reconstructed image of the pin) were examined to simulate the 2D crosssectioning experiment that would be obtained by SEM observation of metallographic sections. For the 3D analysis of the corroded volume, the whole pin was considered. The measured corrosion volumes were similar for both estimations. Nevertheless, the 3D method detects more corrosion in general because small fractions of materials with a high percentage of corroded volume are better resolved. The ability to detect non-uniformly distributed corrosion in samples is a major advantage of the X-ray tomography method compared to conventional 2D methods. 3.3. Corrosion in 2.5 M HCl solution characterized in situ
Fig. 8. AA6111 exposed to 0.7 M HCl for 7 h in situ monitored. Comparison of the corroded volume estimated by the analysis of five cross-sectional slices (2D) and measured in the volume measured in 3D.
To find an explanation for the transition from intergranular corrosion (IGC) to exfoliation-like attack (ELA) and the occurrence of ELA in general, experiments were performed in a more aggressive environment. Based on the assumption that in IGC paths more aggressive solutions build up and based on the observation that ELA initiates at IGC sites, ELA seems to be more pronounced in aggressive solutions. A pin of each alloy was immersed in 2.5 M HCl. The corrosion process was then monitored in situ using a different tomography setting that enables faster data acquisition.
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These settings gave a reasonable balance between signal-to-noise ratio and acquisition time (10 min per scan). Fig. 9 presents a cross-section from the middle of an AA6111 pin immersed in 2.5 M HCl. After an initiation time of about 30 min, rapid localized corrosion showing mainly the directed ELA morphology occurs, but IGC is also visible. IGC apparently did not propagate further between 45 min and 1 h of solution exposure. ELA, however, grew within the ST plane and new ELA sites occurred within this time span. The broadening of the ELA corrosion path is initially fast but stops at a final thickness of about 10 lm. Fig. 10 presents a comparison of the corrosion morphology observed for alloys AA6016 and AA6111 during immersion experiment in 2.5 M HCl. Both alloys show mainly ELA, but as already reported for the less aggressive solution, AA6016 alloy is significantly less susceptible to corrosion. This is derived from a longer initiation time (AA6016 1 h; AA6111 30 min) and the extent of corrosion after 2 h of immersion. Again ELA showed no significant corrosion path broadening for longer immersion times and restricted propagation in ST planes (parallel to sheet surface). These observations imply that like the IGC, the attack may follow a certain microstructural feature. Possible microstructural features that may explain this phenomenon are: (i) grain boundaries and texture, (ii) alignment of the intermetallics and (iii) composition variations in the alloy. In the following of these topics only the influence of grain boundaries and texture
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on ELA is further analysed while a detailed discussion of the ELA mechanism is described in a different paper [42]. 3.4. Exfoliation-like attack (ELA) and grain boundaries/texture Fig. 11a presents a typical EBSD map of a corroded AA6111 (0.7 M HCl, 7 h, ex situ) pin. The different colours represent various grain orientations. The grain boundaries are orientated slightly along the LT direction and no preferential texture is present in the alloy. Fig. 11b shows a SEM image of the identical section. EBSD map in Fig. 11a, revealing most of the grain boundaries, is overlaid on the SEM image. Two sites of ELA, IGC (black) and intermetallics (bright spots) are visible. The IGC follows the grain boundaries, whereas the ELA orientation does not correspond to the grain boundaries. Fig. 12a presents another experimental approach proving the independence of ELA from the grain boundaries in sheet material. It shows a SEM image of a surface of alloy AA6016 exposed to 5 M HCl for 15 min (the increased aggressivity of the solution used for this surface attack followed from the observation that for very acidic conditions the directed ELA strongly dominates the IGC attack). The image was taken from the L plane, looking onto the sheet thickness. It can be seen that the ELA is very distinct and has spacing between the ‘‘lamella” of about 10 lm, which corresponds to
Fig. 9. AA6111 cross-section from the middle position of the pin showing corrosion development as a function of immersion time in 2.5 M HCl solution. IGC is present, and ELA is dominant. One area is magnified to show the IGC and exfoliation-like corrosion together.
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Fig. 10. Comparison of attack in AA6016 and AA6111 after exposure to 2.5 M HCl solution. For both alloys the morphology is similar but the time scale differs.
the ELA path thickness reported previously. The lamella periodicity is significantly smaller than the grain size, and therefore this attack morphology cannot be provoked by the grains or, respectively, the grain boundaries. Fig. 12b shows an EBSD map of the corroded sheet shown in Fig. 12a after polishing away most of the corrosion ‘‘lamella”. The corrosion sites (ELA) visible in black do not follow the grain boundaries. Hence in this experiment, too, ELA did not depend on texture or grain boundaries. 3.5. Dissolution of intermetallic particles within the material
Fig. 11. AA6111 pin corroded in 0.7 M HCl for 7 h (ex situ). (a) EBSD mapping AA6111 showing the grain boundaries and orientation. The pin was polished after corrosion to perform EBSD. (b) SEM/EBSD image; intergranular corrosion (IGC) and exfoliation-like attack (ELA) is visible. Grain boundaries (GB) are marked in white. ELA follows the ST planes.
The behaviour of intermetallics in an aggressive solution was analysed by X-ray microtomography, revealing a new phenomenon. A complete dissolution of micrometer-sized intermetallics, which are reported to be cathodic when characterized at the surface, was observed [43,44]. Fig. 13 illustrates this corrosion attack inside the bulk alloy. The white microscopic features seen within the grey dark matrix are areas of higher absorption coefficient related to the presence of intermetallics. For the alloy presented, these bright features are Al5FeSi and Al15(Fe,Mn)3Si2 intermetallics which also contain a non-stoichiometric amount of copper. Fig. 14 shows the predicted equilibrium composition and intermetallics for the
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4. Discussion Different types of in-depth corrosion propagation phenomena have been successfully investigated by means of X-ray microtomography. The following section discusses different aspects of these types of attack in relation to literature data. 4.1. Surface–deformed layer
Fig. 12. AA6016;exposed to 5 M HCl solution for 15 min: (a) A SEM image of a heavily-corroded surface facing the longitudinal and transverse direction (LT plane). Small aligned lamella (ELA) are visible. (b) EBSD mapping after polishing away most of the lamella (ELA). It can be seen that the exfoliation-like corrosion path is not oriented along the grain boundaries.
alloys used. The thermodynamic calculation model is based on the PANDAT database [45,46]. The black vertical bar shows the final ageing treatment temperature of 180 °C (for 45 min), from which the alloy is quenched to room temperature. For AA6016 (Fig. 14a) mainly Mg2Si, Si and Q(Al5Cu2Mg8Si6) phase will form nanometer-sized [25,47]. At these low homologous temperatures, diffusion is slow and therefore no equilibrium will be reached in 45 min. In addition, primary AlFeSi intermetallics with an average size of about d = 3.5 lm (80% < 4 lm) are present directly from the melt and represent a surface coverage of about 0.8% for these alloys. AA6111 contains primary intermetallics similar in shape and size to those in AA6016 but with a volume content of 1.4%. The main difference from the point of view of corrosion for AA6111 is the non-stoichiometric Cu content in the primary AlFeSi intermetallics, as well as the lower Si content in solid solution. The predicted composition was experimentally verified with electron microprobe (EPMA) measurements. The pin displayed in Fig. 13 was exposed to 0.7 M HCl for 5 h. During this exposure period the Fe containing intermetallics dissolve when an intergranular corrosion path and therefore aggressive solution enriched with different dissolved metallic elements (i.e. MeClx,H+) reaches the intermetallics. After 1 h of exposure, all the intermetallic particles highlighted are still present and no sign of corrosion is detected. After 3 h of immersion, the first corrosion paths intersect with intermetallics. Those intermetallics in contact with the solution leave a black area, indicating that the intermetallics have dissolved. After 5 h of immersion, heavy traces of corrosion at the former intermetallic sites are visible. Hence it may be concluded that cathodic Fe containing intermetallics can dissolve in the inner part of materials during the corrosion process and that local aggressivity can be very high in IGC attack.
At an initial stage, a surface–deformed layer resulting from the machining process is more susceptible to localized corrosion compared to the bulk structure, as reported in literature for submicrometer thick surface layers [27,30,48]. This susceptibility is attributed to the higher defect density present at the surface (e.g. impurities, oxide inclusions, contamination and ultrafine grain structure) [27]. The influence of this deformed layer on filiform corrosion under an organic coating is also discussed in literature [27,49], where it is shown that this type of layer is attacked first and causes further delamination of the coating for certain heat treatments. But a distinction must be made in this case between superficial filiform corrosion and deep penetrating attack. This study was able to show that after the first preferential IGC attack the deformed layer surface does not dissolve further. This indicates a high uniform corrosion resistance in this very aggressive electrolyte. Nevertheless, because this layer is very thin, it is fragile and difficult to detect using conventional methods. This observation raises a question concerning the positive influence of plastic deformation or of the decrease in grain size on corrosion resistance especially for in-depth corrosion propagation. 4.2. ELA vs. IGC and correlation of ELA with grain boundaries To summarize the corrosion initiation and propagation experiments, two types of attack were observed: the IGC and the exfoliation-like attack (ELA). The initiation time is comparable for IGC and ELA whereas ELA develops faster. The results of the experiments performed in the two solutions (0.7 M and 2.5 M HCl) facilitate discussion of the two corrosion phenomena. The assumption that ELA requires higher aggressivity is verified; the apparent aggressivity threshold lies between the two solutions chosen for the experiments. In 2.5 M HCl the two corrosion types compete, but ELA penetrates more deeply (and therefore more quickly) than IGC. In addition, susceptibility of alloy AA6111 to ELA and IGC is higher than that of AA6016. The main differences between AA6016 and AA6111 are the small grain size structure of AA6016 in the center of the pin, higher excess of Si of AA6016 and higher Cu/Mg content of AA6111, which is reflected in a non-stoichiometric Cu content in the intermetallics as well as Cu in solid solution for AA6111. The grain structure is not further considered in this paper as ELA is not seen to depend on the grain boundaries. In AA6016, fine Si-containing intermetallics are present and Si is present in larger amounts in solid solution. Taking these microstructural features into account in relation to the difference between the corrosion susceptibilities of AA6111 and AA6016, it may be concluded that the intermetallics and/or the solid solution content play an important role in the ELA process. The grain boundaries are not responsible for the direction of ELA, especially in aggressive solutions. In a first comparison of intermetallics alignment and ELA corrosion path with the help of X-ray microtomography, no significant correlation was found. Independent of the aggressivity of the bulk solution in all experiments ELA was directed parallel to the former sheet surface (ST plane) and showed a final thickness of 10 lm. These observations are a strong indication that ELA might be correlated to compositional variations in the
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Fig. 13. Sequence of tomograms obtained for AA6111 after 1 h, 3 h and 5 h exposure to 0.7 M HCl. ‘‘Cathodic” intermetallics (white spots) dissolve. After 5 h exposure almost all intermetallics in contact with solution dissolved.
matrix (i.e. Mg, Si and Mn). It is matter of a detailed investigation [42]. 4.3. Dissolution of intermetallics In corrosion experiments where surface processes are considered, the fact that Fe containing intermetallics disappear is mainly attributed to undercutting as a result of dissolution of the surrounding matrix [50], the so called ‘‘trenching process”. It is obvious from Fig. 13 that the corrosion mechanism is different within the bulk of the material, where various solution concentrations induced by the confinement of corrosion products can be found. Trenching around an intermetallic may still be envisaged, but has not been observed in the X-ray microtomography experiments. The complete dissolution of the intermetallics seen can be explained by the fact that if they are not in direct electrical contact with the surrounding Al matrix, they are no longer cathodically protected and can freely corrode. This dissolution can also occur as a result of chemical dissolution due to the extreme aggressivity of the local solution. What has to be stated here is that this is not only a dealloying process often observed for intermetallic particles on the surface and discussed in the introduction, but a complete dissolution. In the case of dealloying, the contrast in microtomography would subsist as it is not the most noble and heavy element that would dissolve. As a result, Fe ions as well as Cu ions are expected to be present in the solu-
tion and can redeposit as further cathodic sites at different locations or act as strong oxidation agents and accelerate the corrosion process. In literature so far it has been stated that the Fe containing intermetallics at the surface undergo selective dissolution of Al, leaving a reactive Fe-, Mn- and Si-enriched remnant. While Fe is said to provoke a higher cathodic activity, Mn and Si are supposed to reduce the cathodic activity because of the formation of an ‘‘inert” passive layer [44]. For AA2024 alloy, extensive work was also performed on Cu-containing S phase particles, where Mg and Al dissolve leaving a free-corroding Cu sponge-like structure with subsequent dissolving Cu redepositing on the surrounding Al matrix [43,51]. Both processes generate greater cathodic activity on the part of the intermetallics in the corrosion process. All the phenomena described do not apply for the complete dissolution of the intermetallics observed, or at least cannot be resolved time-wise if the dissolution is the result of multiple dissolution steps. X-ray microtomography, however, offers an excellent opportunity to detect the complete dissolution of intermetallics within the bulk as a function of time, which is extremely hard or impossible to verify with conventional methods. The location of dissolved intermetallics is detected due to the 3D reconstruction, and the corrosion process can be followed as a function of time within the bulk without taking into account artefacts produced by sectioning or sample preparation. These results raise the question of the chemistry within an intergranular corrosion path.
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path is highly aggressive. It also indicates that the corrosion mechanisms within the bulk may differ from the corrosion attack observed and well-documented at the surface. Acknowledgements The authors would like to thank P. Wyss, Empa Dubendorf, for support in image processing. They are indebted to Axel Stange, Alcan Inc., and Tom Quested, Novelis AG, for help with EBSD measurements and sample preparation. Financial support from the Alcan Technology Fund is gratefully acknowledged. The use of the Advanced Photon Source was supported by the US Department of Energy, Office of Science, Office of Basic Energy Sciences, under Contract No. DE-AC02-06CH11357. References
Fig. 14. Thermodynamic modelling of the equilibrium composition of the alloys with regard to their intermetallic phases for (a) AA6016 and (b) AA6111. The dotted black bar marks the first heat treatment temperature and the black bar at 180 °C marks the final heat treatment temperature.
5. Conclusion In this study the corrosion processes of two AlMgSi alloys were monitored in three dimensions. The synchrotron X-ray microtomography technique was used for in situ characterization. Different issues concerning corrosion morphology and propagation rates were documented, and are discussed: 1. X-ray l-tomography is a useful tool to monitor in situ corrosion processes in the lm-range. Due to the unique possibility of non-destructively viewing inside a large volume of material in <10 min for each scan, time series of corrosion processes characterization within the bulk can be performed and displayed in three dimensions. This facilitates the description of phenomena which are almost impossible to detect directly using conventional methods. 2. Surface–deformed layers are seen to be very susceptible to intergranular corrosion, but after first attack occurs, they are not uniformly corroded further, even though the matrix underneath dissolves completely. 3. A new type of exfoliation-like attack (ELA) was observed experimental with the bulk of the material. The actual mechanism is not yet clear, but the attack does not follow the grain structure, texture or planes of primary intermetallics. ELA was more pronounced for fast corrosion processes in aggressive solution, while for slow processes in less aggressive solutions intergranular corrosion (IGC) prevailed. 4. Intermetallics supposed to be cathodic are seen to dissolve during the corrosion process within the bulk material. This observation provides a hint that the solution within a corrosion
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